High brightness electroluminescent device emitting in the green to ultraviolet spectrum and method of making the same

ABSTRACT

A green-blue to ultraviolet light emitting semiconductor laser having a top contact, a Bragg reflector, cladding layer, active layer, cladding layer, buffer, substrate, bottom contact and a passivation layer. The key aspect is a Ga*N material on a base structure comprising a SiC substrate selected from a group consisting of 2H-SiC, 4H-SiC and a-axis oriented 6H-SiC. Furthermore, the cladding layers have larger band gaps than the active layer and are complimentarily doped.

DESCRIPTION

1. Field of the Invention

The present invention relates to a light emitting optical device capableof emitting light in the green-blue to ultraviolet spectrum, and morespecifically semiconductor devices formed using gallium nitride, e.g.,comprising a gallium nitride active layer on a silicon carbidesubstrate, as well as to a method for forming gallium nitride devicesusing silicon caribide substrates.

2. Description of the Related Art

UV and blue light emitters (both lasers and light emitting diodes(LEDs)) have a wide range of applications including high density opticalstorage, full color displays, color determination systems, andspectroscopic analysis sources. Electroluminescent semiconductor devicessuch as lasers and LEDs that emit light in the green-blue to ultravioletregion of the electromagnetic spectrum are correspondingly of greatinterest, but have not yet reached the levels of performance that arepresently available in red and yellow light emitters, as measured byquantum efficiency and luminous intensity. The major reason for thedeficiency of the blue light-emitting devices is the much less welldeveloped state of the semiconductor materials that are suitable forblue light emission.

Only a few material systems are capable of directly producing solidstate light emission in the blue/UV region of the electromagneticspectrum--the II-VI compounds, silicon carbide, and the III-V nitrides.

Although the II-VI compounds such as ZnSe, ZnS, CdS, and their alloysare direct band gap materials and thus have high optical efficiency,such compounds are plagued by fuidametal material property deficiencies,including softness and low melting temperatures, which cause defects tobe readily generated and propagated, in turn leading to poor reliabilitycharacteristics and short product life. Thus, while LEDs and lasers havebeen demonstrated in II-VI compounds, the lack of stability and theshort lifetimes of the resulting devices are major shortcomings whichhave severely restricted their commercial potential and use. It has beendemonstrated that the short lifetimes in II-VI devices result from therapid propagation of defects throughout their active regions, which actas nonradiative recombination sites. These results are indicative of lowtemperature and room temperature characteristics. Operation at highertemperatures or at high power conditions will increase the degradationrate and exacerbate the problem. Additional problems with the II-VImaterials include the necessity of using quaternary layers to achieveblue emission in relatively lattice-matched heterostructures. Thesimplest and most well developed materials are ZnSe, ZnS, CdS, and theiralloys which produce emission in the blue-green region, about 500 nm.Other II-VI compounds could be used to produce emission in the blue andultraviolet, for example alloys of ZnS and either ZnSe, CdS, CdSe orMnSe. However, these materials have not been well investigated and haveother major problems including the lack of a simple latticematchedheterostructure system and the lack of a suitable substrate.

Blue LEDs formed in silicon carbide epilayers on silicon carbidesubstrates have been reported (U.S. Pat. No. 4,918,497). However,silicon carbide is an indirect band gap material, and thereforeradiative recombination is inefficient, and consequently these SiC-basedlight emitting diodes (LEDs) have poor optical efficiency. Acommercially available 6H-polytype SiC LED with a peak light emission ata wavelength of 470 nm has an external quantum efficiency of 0.02% andperformance of 0.04 lumens/watt (Cree Research, Durham, N.C.). Thisperformance is quite low compared to the best LEDs emitting in the red(AlGaAs, 16% quantum efficiency and 8 lumens/watt) and yellow-green(AlInGaP, 1% quantum efficiency and 6 lumens/watt), where quantumefficiency is defined as the number of photons emitted per electronsupplied X 100%, and luminous intensity is the luminous (visible) fluxoutput of a light source measured in lumens divided by the electricalpower input to the device. Lumens are calculated by multiplying theradiant flux output of a device (in watts) by the human eye'ssensitivity as defmed by the Commission Internationale de L'Eclairage(CIE), and so the luminous efficiency is related to the amount of lightperceived by the human eye per power input.

The performance of SiC LEDs cannot be expected to match that of theIII-V direct band gap materials. In addition, there is no convenientheterostructure system for SiC. A convenient heterostructure system isimportant because it increases device efficiency and permits selectionof the emission wavelength.

The best choice for blue and UV light emitter applications appears to bethe III-V nitrides, i.e. GaN, AlN, InN and their ternary and quaternaryalloys such as AlGaN, InGaN, or AlInGaN. The III-V nitrides meet many ofthe requirements for making light emitting devices. These materialspossess direct band gaps, a convenient, well lattice-matchedheterostructure system, the ability to choose the output wavelength byvarying the composition and structure, and good thermal stability. TheIII-V nitrides exhibit strong luminescence in the ultraviolet and blue.

Gallium nitride (note: as used herein, the term "gallium nitride" isintended to be broadly construed, to encompass gallium nitride per se,as well as gallium nitride alloys compounds, e.g., aluminum, and indium)is a most particularly attractive semiconductor material forelectroluminescent or light emitting devices because its wide band gapprovides for light emission in the green-blue to ultraviolet wavelengthregime. Such light emitting devices require the preparation of galliumnitride layers of high crystalline quality, free of quenching defects.An obstacle to such devices has been the lack of a suitable substrate onwhich to grow good gallium nitride layers. The III-V nitrides and theiralloys are referred to herein as Ga*N materials. As used herein, theterm Ga*N refers to binary (e.g. GaN), ternary (MaN), and quaternary(MM'GaN) type gallium nitride compounds, including, by way of example,such compounds as AlN, InN, AlGaN, InGaN, INAlN, and AlInGaN, wherein Mis a metal which is compatible with Ga and N in the composition MGaN andthe composition MGaN is stable at standard temperature and pressure (25°C. and one atmosphere pressure) conditions, and wherein M' is acompatible metal providing quaternary compounds of the formula M_(1-x-y)M'yGa_(x) N, wherein x and y range from 0 to 1. It will be furtherunderstood that ternary and quaternary compounds may be referred to bygeneral formula without subscripts, e.g., AlGaInN, wherein thestoichiometric coefficients (for aluminum, gallium, and indium, in thisinstance) have been deleted for general reference purposes, it beingunderstood that such alloy compositions entail stoichiometry relative tothe metal components which provides a stable composition at theaforementioned standard temperature and pressure conditions.

Recent improvements in growth and p-type doping have led to severaldemonstrations of high efficiency, GaN-based blue LEDs. GaN LEDs grownon (0001) oriented sapphire have exhibited an external quantumefficiency of 0.18%, almost 10 times that of SiC LEDs (S. Nakamura etal., Jpn. J. of Appl. Phys. 30 (1991) L1998). More recently, AlGaN/GaNheterostructure LEDs with outputs of 1000 mcd at 20 mA (NikkeiElectronics Dec. 20, 1993 (No. 597)) have been reported. By comparison,SiC LEDs emit only about 25 mcd at 20 mA. Additionally, stimulatedemission by photopumping has been demonstrated in GaN (H. Amano et al.,Jap. J. Appl. Phys. 29 (1990) L205); M. A. Khan et al., Appl. Phys.Lett. 58 (1991) 1515). Finally, the III-V nitrides possess manysimilarities to the III-V arsenides GaAs and AlAs, and so growth andfabrication techniques that have been well developed for the lattermaterials may be employed in fabricating nitride-based devices as well.

The band gap of GaN is 3.4 eV while that of AlN is 6.2 eV and InN is2.09 eV. Thus a device with a GaN active layer would emit at about 365nm for band to band recombination. Like GaAs/AlAs, GaN/AlN form aclosely lattice-matched heterostructure system. The difference in thelattice constants of GaN and AlN is about 2.5%. While larger than thedifference between the lattice constants of GaAs and AlAs, this fairlyclose match does permit the use of GaN/AlGaN alloys with mismatches ofless than 0.5%. Such a convenient heterostructure system is importantfor light emitting devices because it increases device efficiency andpermits selection of the emission wavelength. The wavelength can bemodified in several ways. The first is through the use of quantum wellheterostructures, in which the emission energy increases as the wellwidth decreases, because of quantum size effects. The second is the useof AlGaN alloys in the active region, which also increases the emissionenergy relative to GaN. Of course, these two techniques could becombined, if desired. The emission energy can be reduced by the additionof In to the active region alloy.

A key problem with the present GaN light emitting devices is that theyare primarily fabricated on (0001) oriented sapphire (Al₂ O₃)substrates. Use of sapphire has a number of problems. The first is thatthe lattice mismatch between GaN and sapphire is about 13.8%, which isquite large. This large lattice mismatch causes a high density ofdefects at the sapphire/GaN interface, in the range of 10⁸ to 10¹⁰ percm², and these defects propagate up into the device's active regionduring growth of the active layer. Defects pose a serious problem to thereliable operation of optical emitter devices such as lasers and LEDs.Often optical devices have dark line defects which multiply duringoperation. This phenomenon is an especially important problem foremitters grown on lattice mismatched substrates or containing layersthat are mismatched. With continued operation, the density of these darkline defects increases until the light output is reduced to anunacceptable level and device failure occurs. Therefore, the deviceperformance of GaN/sapphire light emitters is limited by crystal qualityeffects.

Another problem with GaN light emitting devices fabricated on sapphiresubstrates is that sapphire is insulating, which means that thestandard, simple LED or laser structures, with one contact on top andthe other on the bottom, cannot be used. Additional fabrication stepsmust be used to make both contacts on the tops of the devices.Additionally, specialized wirebonding and packaging must be employed toaccommodate the non-standard arrangement of the two top contacts. Animportant side effect of having to place both contacts on top of thedevice when a sapphire substrate is used is that the die or chip areamust be about twice as large as the standard structure to leave room forboth contacts on the top surface. This effectively halves the number ofdevices which can be made on a given substrate area, resulting inincreased cost per device.

It would therefore be advantageous to employ substrates other thansapphire for III-V nitride devices.

Silicon carbide has several advantages as a substrate material for III-Vnitride based light emitting devices. The first is that it provides amuch closer lattice match to the III-V nitrides than does sapphire (3.4%vs. 13.8% for sapphire and GaN), leading to devices with fewer defectsand thus higher efficiencies and longer lifetimes. The second is thatSiC can be made conductive, which permits the use of simple andconventional LED and laser structures, with one contact on top and theother on the bottom. Less processing is required thereby, theconventional packages and packaging tools can be used, and the devicerequires about one-half the area of an equivalent device made on asapphire substrate. Finally, most polytypes of silicon carbide show highlight transmittance throughout the visible light wavelengths and someshow good transmittance even into the ultraviolet region. GaN lightemitting diodes that could be fabricated on bulk single crystal, singlepolytype SiC substrates, in particular (0001) oriented 6H-SiC or cubic3C-SiC, have been proposed (U.S. Pat. No. 5,210,051). These combinationsof materials could show significant advantages over the GaN/sapphiresystem. The hexagonal 2H-GaN/(0001)-oriented 6H-SiC system was selectedbecause of the availability of the 6H-SiC substrate and thecompatibility of the crystal structures. The cubic GaN/3C-SiC system wasproposed likewise on the basis of crystal compatibility arguments,although bulk 3C-SiC substrates are not currently available and littleis known of the preparation or properties of cubic GaN. The optimizationof electrical and optical properties, in particular charge carriermobilities and percent transmittance of light, have not been addressed,nor has the use of a heterostructure system, important because itincreases device efficiency and permits selection of the emissionwavelength. Methods to economically fabricate large quantities of GaNLEDs on SiC substrates have not been described. Device structuresdesigned to increase the output of light emitted are needed. Theseissues are crucial to the economical production of high brightness,commercially viable blue light emitting devices. Further, the latticeconstants and thermal expansion coefficients of GaN and SiC aresufficiently different that compatibility problems occur under theconditions of layer growth and subsequent processing.

Differences in the lattice constants and thermal expansion coefficientsof GaN and SiC often cause cracking in GaN layers grown on SiCsubstrates when the samples are cooled from the growth temperature(˜1000° C.) to room temperature. Cracking of the GaN epi-layers is asevere problem for the fabrication of electronic and optical devicesfrom these materials. Consequently, it is very desirable to have amethod to grow GaN layers on SiC that are free of cracks and have verygood structural, optical and electrical properties.

The quality of GaN grown on lattice mismatched substrates such assapphire, Si, GaAs and SiC is greatly improved when a buffer ortransition layer is grown on the substrate prior to growth of the GaNlayer.

Several different types of buffer layers have been used for growth ofGaN on sapphire including AlN (I. Akasaki, H. Amano, Y. Koide, K.Hiramatsu and N. Sawaki, Effects of AlN Buffer Layer on CrystallographicStructure and Electrical and Optical Properties of GaN and Al_(x)Ga_(1-x) N (0<×<0.4) Films Grown on Sapphire Substrate by MOVPE, J.Crystal Growth 98 (1989) 209) and Al_(x) Ga_(1-x) N where 0<×<1 (S.Nakamura, Crystal Growth Method for Gallium Nitride-Based CompoundSemiconductor, U.S. Pat. No.5,290,393.). These buffers are typicallyvery thin (less than 0.5 μm thick, preferably less than 0.1 μm thick)and are grown at low temperature (900° C. or lower). The buffers grownat low temperature are polycrystalline and have a high density ofdefects. The purpose of the polycrystalline buffer is to accommodate thevery large lattice mismatch between sapphire and GaN (13.4%). AlN bufferlayers have also been used for growth on Si substrates (T. Takeuchi, H.Amano, I. Akasaki, A. Watanabe, and K. Manabe, Method of Fabricating aGallium Nitride-Based Semiconductor Device with an Aluminum and NitrigenContaining Intermediate Layer, U.S. Pat. No. 5,389,571).

SiC has a much closer lattice match to GaN (3.4%) than does sapphire.While an AlN layer grown at low temperature has been used as a bufferfor growth of GaN on SiC, improved crystal properties have been obtainedby growing a thin, approximately 1000 Å AlN buffer at temperaturesgreater than 1000° C. (T. W. Weeks, M. D. Bremser, K. S. Ailey, E.Carlson, W. G. Perry and R. F. Davis, GaN Thin Films Deposited viaOrganometallic Vapor Phase Epitaxy on 6H-SiC (0001) usingHigh-Temperature Monocrystalline AlN Buffer Layers, Appl. Phys. Lett. 67(1995) 401.) We have found that improved GaN structural properties, asdetermined by x-ray double crystal rocking curve measurements, can beobtained by using a high temperature AlN buffer. Thefull-width-at-half-maximum (FWHM) for the (0004) GaN reflection is 270arc-sec with an AlN buffer grown at 1100° C. on 6H-SiC. compared to 428arc-secs when AlN grown at 550° C. is used as a buffer on 6H-SiC. GaNgrown on SiC substrates using AlN buffers, however, contains parallelcracks that extend across the entire GaN layer. This is due to strain inthe GaN layer that is caused by the large lattice mismatch between GaNand AlN (2.4%) and AlN and SiC (1.0%). A transition buffer layer hasbeen used to decrease the extent of GaN film cracking when growing onSiC (J. A. Edmond, V. Dmitriev, K. Irvine, Buffer Structure BetweenSilicon Carbide and Gallium Nitride and Resulting SemiconductorDevices", U.S. Pat. No. 5,393,993). This transition buffer contains 3layers which are described as a top (Al,Ga)N layer (%Al>%Ga) anintermediate (Al,Ga)N layer (%Ga>%Al) and a bottom AlN layer. There isstill, however, a large lattice mismatch between the GaN layer and thetop (Al,Ga)N layer of the buffer (>1.4%) in this structure whichintroduces strain into the GaN layer.

In respect of the teachings of U.S. Pat. No. 5,393,993, it is to benoted that growth of GaN on SiC requires the use of specific bufferstructures to achieve high quality material. These typically involvesome graded composition alloys which further complicate that bandstructure. For example, in U.S. Pat. No. 5,393,993 a buffer layer isdisclosed that actually creates a potential barrier to current flowbetween the SiC substrate and the GaN layer. In this structure, the(Al,Ga)N layer next to the device structure has a higher Al compositionthan the (Al,Ga)N layer below it.

These potential barriers lead to increased series resistance and reduceddevice efficiency. Even if the layers are all doped, there will still bethe potential barriers which will cause increased series resistance.This is made clear in the LED structures disclosed in U.S. Patent No.5,393,993. To get around these problems shorting straps need to be addedacross the buffer layer. This greatly complicates the fabricationprocess and results in increased costs. While applicable for LEDs, withlow current densities, this solution is unsuitable for high currentdensity devices such as lasers and high power electronic devices.

Another problem when growing GaN on foreign substrates is that thegrowth conditions need to be very carefully controlled during the intialnucleation stages to achieve high quality material. Typically thinbuffer layers (about 200 Å) are grown at low temperature (about 500° C.)prior to the growth of the GaN device structure. The thickness, growthrate and growth temperature must be controlled to very tight tolerancesto achieve good material.

The buffer layer described hereinafter addresses both of these problems.In the first place, these can be grown at the GaN growth temperature,thus eliminating the need for control at two temperatures in the growthprocess, as well as provide tight control over the thickness of the lowtemperature buffer layer used in conventional processes.

The conductive graded buffer starts with a low Al composition and gradesdown to GaN. Thus the bandgap discontinuity is much less, resulting in amuch lower barrier to current flow, and thus a lower series resistanceand higher device efficiency. In the disclosed buffer layer there is amuch smaller barrier to current flow and no potential wells to trapcarriers. These improvements result in lower series resistance andhigher device efficiency in devices where the substrate comprises onecontact to the device.

It is therefore an object of the present invention to provide brightgreen-blue to ultraviolet light emitting devices with high opticalefficiency, long lifetime, and a simple fabrication process which iscompatible with current device fabrication, testing and packagingprocesses. The invention provides improved charge carrier mobilities andlight transmittance. Various objectives of the invention includeproviding: a method for making these bright blue light emitting devices;a method for selection of the output wavelength of the light emittingdevice by control of the composition and structure of the active regionof the device; a device structure that enhances the efficiency of lightextraction from the light emitting device by means of Bragg mirrors thatreflect light of the wavelength emitted by the device; an alternativemethod for device fabrication, to keep damage from the dicing operation,which may degrade the device characteristics, away from the activeregion and thereby avoid the necessity of post-dicing etching steps;methods for reducing the misfit dislocation density at the interfacebetween the substrate and the active layer. Another object of thepresent invention is to provide composite structures comprising bufferlayers between high-quality, crack-free layers of gallium nitride onlattice-mismatched substrates, especially silicon carbide, and a methodfor preparing such composite structures.

SUMMARY OF THE INVENTION

In one aspect, the present invention relates to a brightgreen-blue-to-ultraviolet light emitting optical device, e.g. agreen-blue to ultraviolet emitting laser or a green-blue to ultravioletemitting diode, comprising a Ga*N active layer and a silicon carbidesubstrate of an anisotropic polytype such as hexagonal or rhombohedral,preferably oriented such that growth of the Ga*N layers takes place onor slightly off-axis on an a-axis or c-axis oriented substrate. The SiCsubstrate preferably consists of 2H-SiC, 4H-SiC, or a-axis oriented6H-SiC.

In another aspect, the invention may incorporate a structuralmodification to increase the light o utput comprising a dielectric Braggmirror beneath the LED structure, made of alternating layers of AlN,GaN, InN or their alloys.

In yet another aspect, the devi ce may be defined during devicefabrication by a process comprising mesa etching, optional passivationof the mesa edge, and contact formation, to produce a device in whichthe critical p-n junction edge region is physically separated from thedicing region.

Another asp ect of the invention relate s to a semiconductor device orprecursor thereof. comprising a layer of single crystal silicon carbideand a layer of single crystal gallium nitride, having a buffer layertherebetween comprising a compositionally graded Ga*N layer.

Such compositionally graded Ga*N layer may for example comprise acompositionally graded Al_(x) Ga_(1-x) N buffer layer between thegallium nitride and silicon carbide layers, wherein x can range from 0to 1, and the buffer layer is compositionally graded from an interfaceof the the buffer layer with the silicon carbide layer at which x is 0,to an interface of the buffer layer with the gallium nitride layer atwhich x is 1.

A still further aspect of the invention relates to aspect of theinvention relates to a semiconductor device or precursor thereof,comprising a silicon carbide substrate and an epitaxial layer of galliumnitride, having a buffer layer therebetween comprising a compositionallygraded (Al,Ga)N buffer layer. A preferred silicon carbide substratecomprises 6H-SiC, with an Al_(x) Ga_(1-x) N buffer layer where the Alcomposition (x) is graded from 1 at the buffer-SiC interface to 0 at theGaN-buffer interface. The thickness of such graded buffer can range from200 Å up to 5 μm, with a preferred range being from 500 Å up to 1 μm.

A variation of the above-described structure may comprise a thin AlNbuffer layer which is initially grown on SiC followed by thecompositionally graded (Al, Ga)N buffer and the GaN epitaxial layer. Thethickness of the AlN buffer layer can range from 50 Å up to 5 μm, with apreferred range being from 100 Å up to 1 μm.

Another aspect of the invention relates to aspect of the inventionrelates to a semiconductor device or precursor thereof, comprising aconductive buffer layer of the formula Al_(x) Ga_(1-x) N wherein thebuffer layer includes a lower Al composition layer that is doped torender it highly conductive in character, such that x is in the range offrom x=1 to x=0.3, preferably with x being between 0.7 to 0.2. Thecorresponding end composition could be GaN or a lower Al_(y) Ga_(1-y) Ncomposition, where y<x. The use of the composition Al_(y) Ga_(1-y) Ninstead of GaN as the final composition in the grade may be particularlyuseful when the epitaxial layer to be grown on the buffer is (Al, Ga)Ninstead of GaN.

The conductive buffer layer may be made conductive by doping with anysuitable dopant species, so that the layer is of the desired n-type orp-type, e.g., silicon or magnesium, respectively.

The Ga*N buffer layers described hereinabove may be made of (Al, Ga)Ncompositionally graded layer, (Al, In)N compositionally graded layer or(Al, Ga, In)N compositionally graded layers.

If an epitaxial layer to be grown on the buffer contains In, thecompositionally graded buffer layer may be continued to the (In, Ga)N,(In, Ga)N, or (Al, Ga, In)N composition of the epitaxial layer. Forexample, if the epitaxial layer is In₀.1 Ga₀.9 N, then the bufferstructure may illustratively begin with AlN and grade through (Al, Ga)Nto GaN, and then continue grading from GaN to In₀.1 Ga₀.9 N.

Another aspect of the invention relates to a semiconductor device orprecursor thereof, comprising a SiC substrate element including amesa-shaped portion having a Ga*N quantum well structure or other Ga*Nstructure formed thereon, wherein the mesa-shaped portion has anequivalent diameter in the range of from 50 to 300 μm, and a height offrom 1 to 15 μm The SiC substrate may comprise 2H-SiC, 4H-SiC, or a-axisoriented 6H-SiC. The quantum well structure may have a well width offrom 150 Å to 750 Å.

Other aspects, features and embodiments of the invention will be morefully apparent from the ensuing disclosure and appended claims.

DESCRIPTION OF THE DRAWINGS

FIG. 1 shows a schematic representation of a hexagonal silicon carbidecrystal with the vertical arrow indicating the (0001) direction; the"a-axis" directions are denoted by the horizontal arrows, and the c-axisdirection by the vertical arrow.

FIG. 2 shows percent transmission of light in the wavelength range 200to 800 nanometers as a function of wavelength for 4H-SiC.

FIG. 3 shows percent transmission of light in the wavelength range 200to 800 nanometers as a function of wavelength for 6H-SiC.

FIG. 4 shows a schematic of a light emitting diode according to thepresent invention, comprising a conductive SiC substrate, a Ga*N activeregion, a heavily doped Ga*N contact layer, and top and bottomelectrical contacts.

FIG. 5 shows a schematic of a light emitting diode according to thepresent invention, comprising a conductive n-type 4H-SiC substrate, ann-type AlGaN barrier layer, a GaN active region, a p-type AlGaN layer,and a heavily doped p-type GaN contact lay incorporating a p-typecontact to the top and an n-type contact to the bottom.

FIG. 6 shows another embodient of an LED according to the presentinvention, incorporating a structural modification to increase the lightoutput comprising a dielectric Bragg mirror beneath the LED structure,made of alternating layers of AlN, GaN, InN or their alloys.

FIG. 7 shows a device which is defined during device fabrication by aprocess comprising mesa etching, passivation of the mesa edge andcontact formation, to produce a device in which the critical p-njunction edge region is removed from the dicing region.

FIG. 8 shows a GaN/SiC composite including a GaN layer grown on acompositionally graded (Al, Ga)N buffer layer on a SiC layer.

FIG. 9 shows a variation of the structure shown in FIG. 8, in which theGaN/SiC composite includes a thin AlN buffer layer is initially grown ona SiC substrate followed by the compositionally graded (Al, Ga)N bufferlayer and a GaN epi-layer.

FIG. 10 shows an individual mesa suitable for fabricating asemiconductor device or precursor thereaccording to the invention.

FIGS. 11, 12, and 13 show a schematic representation of an exemplaryprocess used to produce mesa-defined devices in accordance with theinvention.

FIG. 14 shows a schematic representation of a semiconductor laseraccording to one embodiment of the present invention.

FIG. 15 depicts,another example of a laser structure according to thepresent invention.

FIG. 16 shows a schematic sectional elevation view of a light emittingdiode fabricated with a gallium nitride active layer, grown on acompositionally graded (Al, Ga)N buffer layer, on a silicon carbidesubstrate.

FIG. 17 is a plot of intensity, in arbitrary units, as a function ofwavelength, for GaN band-edge PL intensity from material grown on top ofthe mesas (-) and on off-mesa regions ( . . . ) of an SiC substrate.

FIG. 18 is a plot of intensity, in arbitrary units, as a function ofwavelength, for a 750 Å GaN quantum well structure grown on top of themesas (-) and on off-mesa regions ( . . . ) of an SiC substrate.

FIG. 19 is a plot of intensity, in arbitrary units, as a function ofwavelength, for a 150 Å GaN quantum well structure grown on top of themesas (-) and on off-mesa regions ( . . . ) of an SiC substrate.

FIG. 20 is a graph of flow rate as a function of time, showing thevariation of TMAl, TMGa flows during growth of a buffer.

FIG. 21 is a graph of intensity, in arbitrary units, showing a doublecrystal x-ray rocking curve of a GaN layer grown on a compositionallygraded (Al, Ga)N buffer on 6H-SiC.

FIG. 22 is a graph of simulated and experimentally measured reflectivityfrom a 24 period 42.8 nm GaN/44.6 nm Al₀.15 Ga₀.85 N reflector stack.

FIG. 23 is a graph of PL spectra as a function of input pump power for aGaN/AlGaN Bragg reflector stack grown on a SiC substrate.

FIG. 24 is a graph of PL spectra as a function of input pump power for aGaN/AlGaN Bragg reflector stack grown on a sapphire substrate.

FIG. 25 is a schematic cross-section of VCSEL structures of a device inwhich the Bragg reflectors are 30-period Al₀.40 Ga₀.60 N/Al₀.12 Ga₀.88 N(397 Å/372 Å) multilayer stacks.

FIG. 26 is a double crystal x-ray rocking curves (w-scan) of VCSEL grownon (A) 6H-SiC and (B) sapphire substrates.

FIG. 27 is a graph of room temperature surface emission spectra forVCSEL grown on sapphire at increasing pump intensities.

FIG. 28 is a graph of peak and integrated intensities of theVCSEL/sapphire emission spectrum as a function of pump intensity inwhich the extrapolated threshold pump intensity is 2.0 M W/cm².

FIG. 29 is a graph of room temperature surface emission spectra forVCSEL grown on SiC at increasing pump intensities.

DETAILED DESCRIPTION OF THE INVENTION, AND PREFERRED EMBODIMENTS THEREOF

The present invention is based on the discovery that by using siliconcarbide of a selected polytype and growing Ga*N on that substrate inselected orientations, light emitting devices, e.g. a green-blue toultraviolet emitting laser or a green-blue to ultraviolet emittingdiode, of surprisingly high quantum efficiency and luminous efficiencymay be fabricated. The luminescence efficiency, charge carriermobilities and transmittance of light are improved.

The substrate polytype and orientation are selected to optimize theperformance of the light emitters. The properties that figure moststrongly in this selection are crystal structure, charge carriermobilities, and transparency of the substrate to light of thewavelengths desirably emitted by the light emitters.

SiC has over 200 polytypes, a situation that is a consequence of aone-dimensional variation in the stacking sequence of the Si-C layers.The polytype with the largest band gap is the 2H hexagonal form, with aband gap of 3.3 eV, which corresponds to a band-to-band emissionwavelength of 380 nm. While the ensuing discussion is primarily directedto the 2H-SiC, 4H-SiC and 6H-SiC polytypes, it will be recognized that awide variety of SiC polytypes may be utilized in the broad practice ofthe present invention. Generally preferred polytype species will havemore isotropic crystal structures, high carrier mobilities, and lowanisotropy of mobility.

GaN grows in hexagonal (2H) and cubic polytypes. The hexagonal polytypesgrow on hexagonal SiC polytypes, and cubic GaN grows on 3C-SiC. Highcrystal quality hexagonal GaN epilayers can be obtained on 2H-, 4H- and6H-SiC substrates.

In silicon carbide, the mobility of the charge carriers varies with thepolytype. To maximize device performance, specific orientations andpolytypes of SiC are selected to achieve a very high useful mobility.The mobility of 4H-SiC is about twice that of 6H-SiC, and that of 2H iseven higher yet. Thus the use of 4H-SiC or 2H-SiC substrates immediatelyprovides an advantage when a highly conductive substrate is desired, asit is for LED and laser devices. Conductivity s behaves as s=q μn whereq is the electronic charge, μ is the carrier mobility and n is thecarrier concentration. Because the mobility is twice as high in 4H-SiCcompared to 6H-SiC, the conductivity can be made twice as high.

Another method to achieve high useful mobility is to ensure that thecrystal's high mobility axis is aligned with the current flow direction.FIG. 1 shows a schematic of hexagonal SiC crystal 10 showing thecrystallographic directions and sample orientations for measuringcarrier mobility parallel to (μ//) and perpendicular (μ⊥) to the (0001)face. The mobilities in 6H-SiC measured in the (0001), 1010 and 1120directions are given in Table I below:

                  TABLE I    ______________________________________    Carrier Concentration                     Orientation                                Mobility    ______________________________________    2 × 10.sup.18                     1010 ("a-axis")                                104    2 × 10.sup.18                     1120 ("a-axis")                                79    2 × 10.sup.18                     0001 ("c-axis")                                35    2 × 10.sup.18                     0001 ("c-axis")                                28    ______________________________________

It is clear that the mobility in the (0001) direction is much lower thanthat in the a-axis directions. Thus for the polytypes of SiC which havemore anisotropic crystal structures, the use of specific orientationsprovides higher mobility and thus higher conductivity for deviceapplications. Since the crystal structure is less anisotropic for 2H and4H-SiC than for 6H-SiC, the mobility anisotropy is much less than in6H-SiC. This approach is especially pertinent in the case of 6H-SiCwhich is the polytype that is the most readily available commercially,and, even though it does not have best properties, may be the practicalselection for substrates.

To summarize, two methods are presented to achieve higher conductivitiesin SiC substrates and thus less power loss and more efficient operationin optical emitters. These are (1) to use higher mobility SiC polytypes,i.e. 2H and 4H-SiC, for the substrate and (2) to use specificorientations of anisotropic crystals to align the high mobility axiswith the current flow direction, i.e. the a-axis orientation of the6H-SiC polytype.

The high carrier mobility has two positive impacts on deviceperformance. First, in the active region of the device, the mobility inthe current flow direction is large, which results in improved deviceperformance. Second, the substrate is a parasitic resistance in serieswith the device, and therefore, a high mobility will result in a lowsubstrate resistance. Small substrate resistance means a lower forwardvoltage drop and less power consumed (wasted) in the substrate, i.e. amore efficient device. In addition, less heat will be generated as aresult of the substrate resistance. It should be notes that SiC has athermal conductivity is about 10 times larger than that of sapphire.This higher thermal conductivity aids in conducting heat away from theactive junction region.

Optical properties of the substrate are also important. In this regard,another significant advantage of the 4H-SiC and 2H-SiC substratesrelative to 6H-SiC is their larger band gaps. The larger band gaps makethese substrates effectively transparent for shorter wavelength lightemission. This means that 2H- and 4H-SiC substrates are effectively moretransparent over a wider wavelength range than 6H- or 15R-SiC. Since inLEDs a great deal of light can come from out of the substrate, theselection of 2H or 4H polytypes offers a significant advantage. FIG. 2shows a graph of the percent transmission of light in the wavelengthrange 200 to 800 nanometers as a function of wavelength for 4H-SiC andFIG. 3 shows the corresponding graph of the percent transmission oflight in the wavelength range 200 to 800 nanometers as a function ofwavelength for 6H-SiC. It is clear that 4H is transparent over a widerwavelength range. Although the relative percent transmission numbers arenot directly comparable because the percent transmission above the bandedge wavelength also depends on the doping density, the wavelength atwhich transmittance goes to zero, or "cut-off" wavelength, isindependent of the dopant concentration.

As FIGS. 2 and 3 shows, the cut-off wavelength is significantly lowerfor the wider band gap 4H-polytype than for 6H-SiC, making thissubstrate suitable for light emitting devices that emit at higher energyultraviolet wavelengths. The high energy transmission of 4H- and 2H-SiCenables higher efficiency LEDs that emit in the green-blue toultraviolet range. High energy LEDs have uses in communicationsapplications, where they can carry a higher density of information, and,in spectroscopy applications, where they can excite a wider range ofenergy level transitions.

As the concentration of dopants increases, transmittance decreases, andso conductivity and transmittance can be balanced in a trade-offsituation. However, the higher carrier mobilities of the 2H-SiC, 4H-SiCand a-axis oriented 6H-SiC combined with high transmittance at shortwavelengths allow doping levels to be selected for a device withoptimized conductivity and transmittance, which has good electrical andoptical properties.

Therefore, to summarize, substrates comprising hexagonal 2H-, 4H- anda-axis oriented 6H-SiC polytypes are preferred, and of the hexagonalpolytypes, the 2H-SiC and 4H-SiC polytypes, which have wider band gapsthan the 6H-SiC polytype, as well as higher carrier mobilities, are mostpreferred.

In contrast to the use of SiC as a substrate for GaN growth, manydifferent types of devices have been made in SiC. Virtually all of thesedevices have been made on c-axis (0001) oriented 6H-SiC substrates. Forall hexagonal SiC polytypes, there are actually two specific low indexdirections perpendicular to the c-axis: (1010) and (1120), which aregenerically called "a-axis." To date only one device has been proposedto be made on an a-axis oriented substrate. This is in an IMPATTdevelopment program at Westinghouse, funded by the Naval Weapons Centerat China Lake, Calif. ("High Power Silicon Carbide IMPATT DiodeDevelopment," Eldridge et al., 2nd Annual AIAA SDIO InterceptorTechnology Conference, Jun. 6-9, 1993, Albuquerque, N.M.). This IMPATTdevice is an electronic, not an optical device. The a-axis was chosen inthis prior art work to achieve certain ionization rates when the deviceis operated in avalanche (reverse bias) conditions, which are totallydifferent from the reasons for using a-axis orientation for LED or laserdevices. Additionally, the device operating conditions are totallydifferent (LED and laser devices operate in small forward biasconditions). To date, no actual devices have been reported to have beenmade on a-axis oriented substrates. GaN LEDs and lasers made on a-axisoriented SiC provide many advantages that have evidently not beenpreviously appreciated.

While LEDs made using 6H-SiC epilayers on 6H-SiC substrates have beenreported previously (U.S. Pat. No. 4,918,497), and GaN LEDs weredescribed on c-axis (0001) oriented 6H-SiC substrates (U.S. Pat. No.5,210,051), to date, there have been no reports of GaN-based opticaldevices (LEDs or lasers) grown on 2H-SiC, 4H-SiC or a-axis 6H-SiCsubstrates. See Y. Ueda, T. Nakata, K. Koga, Y. Matsushita, Y. Fujikawa,T. Uetani, T. Yamaguchi and T. Niina, Mat. Res. Soc. Symp. Proc. vol.162, 1990 Materials Research Society, "Liquid Phase Homoepitaxial Growthof 4H-SiC Crystals and fabrication techniques of bluish-purple lightemitting diodes."

Moreover, in combination with epilayers of the AlGaInN alloys, whosewavelength of light emission can be selected by modifying compositionand thickness, these wider band gap polytypes show distinct andpreviously unappreciated advantages.

When the advantages of the electrical and optical properties of 2H-, 4H-and a-axis oriented 6H-SiC substrate s are combined with growth of Ga*N-based active regions, with their high optical efficiency, a new highpower optical emitter device structure is produced. GaN and relatedmaterials are very efficient optically because of their direct band gap.In addition, the ability to make heterostructures with alloys of AlN andInN enables heterostructure devices as well as the ability to select theoutput wavelength by choice of composition and/or thickness of theactive region.

FIG. 4 is a schematic sectional elevation view of a generalized LEDstructure 12 according to the present invention. The LED structure 12comprises a green-blue to ultraviolet light emitting Ga*N material 13 ona base structure 14 comprising a SiC substrate. The diode structure 12includes a p-n junction comprising Ga*N layers 13 and 17, a contact 15on the upper sur f ace of the Ga*N layer 17, and a contact 16 on thebottom surface of the base structure 14. The substrate is selected fromany orientation of 2H-SiC or 4H-SiC, or a-axis oriented 6H-SiC.

A schematic of an electroluminescent device according to the inventionis shown in FIG. 5. The device 20 has a conductive substrate 26. Thesubstrate may be n-type or p-type. On the substrate is formed a Ga*Nbarrier layer 25, a Ga*N active region 24 and a Ga*N barrier layer 23.Finally the structure is completed with a heavily doped contact layer22, wherein "heavily doped" means a doping level of ≧10¹⁸ cm⁻³, andpreferably ≧10¹⁹ cm⁻³. Contact 21 is formed to heavily doped contactlayer 22 and contact 27 is formed to the conductive SiC substrate 26 onthe lower sides of the device. The Ga*N layers 23 and 25 have a largerband gap than the Ga*N active region, which confines the injectedcarriers and prevents them from diffusing out of the active region. Thisresults in greater optical efficiency. The contact layer 22 should haveas small a band gap as possible, making top contact formation easier.The device is fabricated by making a contact to the top and a contact tothe bottom. The wafer is then diced up into individual devices.

In one embodiment of the device of the general type shown in FIG. 5, thedevice 20 has a conductive, n-type 4H-SiC substrate 26. On thissubstrate is formed a n-type AlGaN barrier layer 25, a GaN active region24 and a p-type AlGaN layer 23. Finally the structure is completed witha heavily doped p-type GaN contact layer 22. Contact 21 is formed to theheavily doped p-type GaN contact layer and contact 27 is formed to then-type SiC substrate on the lower sides of the device. The AlGaN layershave a larger band gap than the GaN active region, which confine theinjected carriers and prevents them from diffusing out of the activeregion, leading to greater optical efficiency. The p-type GaN contactlayer has a smaller band gap than the AlGaN, making top contactformation easier. The device is fabricated by making a p-type contact tothe top and an n-type contact to the bottom. The wafer is the n diced upinto individu al devices.

Doping the Ga*N layers p-type can be achieved using Mg, C, or Zn as thedopant, introduced during growth, with Mg preferred. Source reagentssuch as CCl₄ (carbon dopant), bis(cyclopentadienyl)magnesium (Mgdopant), diethylzinc or dimethylzinc (Zn dopant), etc. may be employedto supply the dopants during deposition. P-Type dopants can be activatedby thermally annealing in a non-hydrogen-containing ambient (Ar, N₂, butnot NH₃) or by LEEBI (low energy electron beam irradiation). However,the latter technique works only to the depth of the beam penetration,e.g. only a few microns using energies that are easily accessible in anelectron microscope source.

Doping of the Ga*N layers n-type can be achieved in situ using Si, Se,or Te as dopant, introduced during growth, with Si preferred. Silane,disilane, hydrogen selenide or telluride, or organometallic Si, Se, orTe compounds are suitable precursors for the introduction of thesedopants during growth. Contact materials that are suitable for the topcontact include Ti/Mo/Au, Au/Ni or Au/Ge/Ni alloys or Au metal. For thebottom contact to the silicon carbide substrate, Ni. Pd, W or Ta may beemployed.

It should be understood that the compositions of the barrier and activeregions of the device shown in FIG. 5 may be modified to achievedifferent performance. The major aspect determining the composition ofthe active region is the desired output wavelength. This is determinedby the band gap of the active region. The band gaps vs. latticeconstants for a number of materials including the III-V nitrides aretabulated in Table II below.

                  TABLE II    ______________________________________    Material      Band Gap Lattice Constant    ______________________________________    AlN           6.28     3.112    GaN           3.45     3.16    InN           2.09     3.544    AlP           2.45     5.467    GaP           2.26     5.4512    InP           1.35     5.8686    AlAs          2.15     5.6605    GaAs          1.42     5.65325    InAs          0.36     6.0584    ZnS           3.68     5.42    CdS           2.42     5.8320    CdSe          1.70     6.050    CdTe          1.56     6.482    Si            1.12     5.43095    3C-SiC        2.2      4.36    6H-SiC        2.9      3.09    ZnO           3.35     4.580    Sapphire      --       4.758    ______________________________________

It is clear that adding AlN to GaN increases the band gap while addingInN to GaN decreases the band gap. The band gap of GaN is 3.4 eV whichcorresponds to a wavelength for band-to-band light emission of about 365nm. For a blue LED, the wavelength should be in the range of 440 to 480nm, which requires the addition of 5-50 mole % In to a GaN activeregion. In other words, the active region should have a composition ofIn_(x) Ga_(1-x) N where 0.05<x<0.50 to produce a blue LED. This rangewill produce different shades of blue. This wide range is specifiedbecause the emission may come from band-to-band or impurityrecombination. If it is band-to-band, then the output wavelength isclearly determined by the band gap. On the other hand, if it is impurityrelated recombination, then the output wavelength is less than theenergy of the band gap by approximately the ionization energies of thedopants. The structure depicted in FIG. 3 can be used to provide lightemission from the UV to visible by appropriate variation of thecompositions in the active layer 24.

In another embodiment of an LED emitting in the blue region of theoptical spectrum, x in the composition In_(x) Ga_(1-x) N ranges from0.05<x<0.50, preferably in the range of 0.07<x<0.15. In the Al_(y)Ga_(1-y) N barrier layers, y could vary from 0 to 0.3, preferably 0 to0.15.

An alternative version would use a p-type substrate, a p-type barrierlayer followed by the active region and an n-type barrier layer andn-type contact layer.

It is also understood that the lower band gap contact layer on the topof the structure may be made of an InGaN alloy or InN for improvedcontacts, or alternatively the contact layer may be omitted.

Other versions of this device utilizing the conductive SiC substratecould have active regions made out of quantum wells. In this case, theactive region is thin enough that quantum size effects as well as layercomposition control the emission wavelength. in this case the emissionwavelength is determined by the well width, the well composition and thebarrier composition. In general, as the well width is decreased, theemission energy is increased.

The active region may also consist of a Ga*N material of lower band gapthan the cladding layers 23 and 25, but instead of having an activeregion of uniform composition, could have one or more wells of Ga*Nlayers of even lower band gap material than the bulk of the activeregion. The carriers will thermalize into those wells and emit lightwith a wavelength dependent on the band gap and thickness of the wellmaterials.

An additional structural modification to increase the light outputemploys a dielectric Bragg mirror underneath the LED structure. TheBragg mirror structure may suitably comprise sequential layers ofmetallonitride materials (compounds or alloys), e.g. alternating layersof AlN, GaN, InN or their alloys, "Ga*N," where the term Ga*N refers tobinary (e.g. GaN), ternary (MGaN), and quaternary (MM'GaN) type galliumnitride compounds, including, by way of example, such compounds as AlN,InN, AlGaN, InGaN, InAlN and AlInGaN, wherein M is a metal which iscompatible with Ga and N in the composition MGaN and the compositionMGaN is stable at standard temperature and pressure (25° C. and oneatmosphere pressure) conditions, and wherein M' is a compatible metalproviding quaternary compounds of the formula M_(1-x-y) M'yGaxN, whereinx and y range from 0 to 1.

This Bragg mirror structure acts as a mirror for specific wavelengths,and increases the light output by efficiently reflecting light up andout of the device. FIG. 6 shows a schematic of this type of structure,where the light emitting device 30 comprises a silicon carbide substrate37, wherein the silicon carbide substrate is preferably of the 2H, 4H,or a-axis 6H polytype, upon which is grown a Bragg mirror 36 comprisingalternating layers of AlN, GaN, InN or their alloys. A barrier layer 35of n-type Ga*N is grown on the Bragg mirror, followed by the activeregion 34 which may comprise Ga*N, and a p-type Ga*N layer 33. Finallythe structure is completed with a heavily doped p-type Ga*N contactlayer 32, wherein "heavily doped" means a doping level of ≧10¹⁸ cm⁻³,and preferably ≧10¹⁹ cm⁻³. Contact 31 is formed to heavily doped p-typeGa*N contact layer and contact 38 is formed to the SiC substrate on thelower sides of the device. As described previously, the Ga*N layers 33and 35 have band gaps larger than the active region 34. The contactlayer has as small a band gap as possible or may be omitted. Of course,the doping scheme could be reversed, using a p-type substrate, and soforth.

The Bragg reflector acts as a frequency-dependent mirror, that is thepeak in the reflectivity spectrum is determined by the layer thicknessesin the structure. The design of the Bragg reflector is governed by thedielectric reflector expression (W. Driscoll and W. Vaughan eds,"Handbook of Optics," McGraw-Hill, N.Y., 1978, Chp. 8.):

    n.sub.1 t.sub.1 +n.sub.2 t.sub.2 =1/2

where n₁ and n₂ are the indices of refraction at the wavelength 1 and t₁and t₂ are the thicknesses of the materials in the reflector. The totalreflection from the reflector is given by:

    Rmax={ (n.sub.m /n.sub.s)-(n.sub.1 /n.sub.2).sup.2N !/ (n.sub.m /n.sub.s) +(n.sub.1 /n.sub.2).sup.2N !}.sup.2

where N is the total number of periods of materials 1 and 2 and n_(i)t_(i) for each of these layers is an odd multiple of 1/4. The indexn_(m) is for the medium into which the light is reflected and n_(s) isthe index of the substrate on which the stack resides. This expressiongives the peak reflectivity; the peak reflectivity decreases as thewavelength varies either higher or lower than 1.

The refractive index of AlN is reported to be n_(AlN) ≈2.1. Therefractive index of GaN in the blue part of the spectrum lies nearn_(GaN) 2.6-2.7. Using these values along with the wavelength of maximumemission for a GaN diode of ≈430 nm, the thicknesses of the GaN and AlNlayers are 41 and 51 nm, respectively (assuming maximum reflection fromthe mirror). It should be noted that the index ratio for GaN/AlN is verysimilar to that of GaAs/AlAs, which indicates that highly reflectivemirrors will be achievable.

To date there has been only one report of a Bragg reflector made in theAlGaN system. This utilized 18 periods of GaN and Al₀.2 Ga₀.8 N layers.The peak reflectivity was measured to be 80% at 442nm for this structure(M. A. Kahn, J. N. Kuznia, J. M. Van Hove, D. T. Olson, Appl. Phys.Lett. 59, 1449 (1991)).

In addition to reflectivity, the resistance of the Bragg mirror is alsovery important. Current must flow through the reflector, and if it has ahigh resistance, thermal heating and device degradation occur. Even ifdevice destruction does not occur, overall device efficiency willsuffer. One problem with the simple Bragg reflector is that thepotential barriers in the superlattice (which result from different bandgaps of these materials) impede the carrier flow and result in largeseries resistances.

There are a number of methods for reducing the series resistance of theBragg reflector. The first is to make the material in the reflector asconductive as possible. However this does not reduce the potentialbarriers in the conduction and valence band. These barriers can bereduced by grading the composition at the interface instead of having anabrupt change. Grading can be done in two ways. The first is to changethe flux of the constituent species as a function of time. The second isto insert either a layer of intermediate composition or a short periodsuperlattice at the interface between the two main layers.

LEDs are made in the simplest possible way to achieve low costs and highyield.

Many LEDs are defined simply by the dicing operation that occurs toseparate them. Damage induced at the edge of the LED during the dicingoperation may cause a decrease in the radiative efficiency or lifetimeof the device. Such damage is often removed by performing an etch afterdicing to remove the damaged material.

The present invention encompasses an alternative to such etching, whichinvolves fabrication of the individual devices by a mesa etchingtechnique which provides an alternative structure which does not requirethis etching. Device definition may thus be performed during thefabrication step to keep damage from the dicing operation away from theactive region.

The conventional LED structures are shown in FIGS. 4 and 5. Thesedevices are defined by the dicing operation. FIG. 7 shows a device whichis defined during device fabrication by a process consisting of mesaetching, passivation of the mesa edge and contact formation. Finally thewafer is diced to separate the devices. The major difference betweenthis structure and the one shown in FIG. 5 is that in the mesa device,the critical p-n junction edge region is spatially separated from thedicing region. This prevents any dicing damage from encroaching into theactive region, which may degrade the device characteristics.

In FIG. 7, the LED 40 comprises a silicon carbide substrate 46 followedby a n-type Ga*N barrier or cladding layer 45, a Ga*N active layer 44,and a p-type cladding Ga*N layer 43, and a heavily doped p-type Ga*Ncontact layer 42. Contact 41 is formed to heavily doped p-type GaNcontact layer and contact 47 is formed to the SiC substrate on the lowersides of the device. A passivation layer 48 protects the device fromdegradation and reduces surface leakage currents. The passivation layermay comprise silicon dioxide, silicon nitride or other materials. Layersof these materials are deposited using chemical vapor deposition,sputtering, plasma-assisted deposition, or other layer-forming processesknown in the art. The passivation layer thickness may range from 200 to2000 Å. The cladding layers 43 and 45 have larger band gaps than theactive layer 44. These devices are made by masking of the grown LEDstructure and etching the mesas after growth.

Typical metallurgy for the p-type contact uses aluminum or gold/aluminumalloys which are annealed at temperatures from 700° C. to 1000° C.N-type contact metallurgy includes Ni, W, Pd, and Ti. These metals areannealed at 800° C. to 1100° C. to form the contact, excluding Ti whichis ohmic as deposited.

The wafer is then sawed using a dicing saw to produce individual LEDdevices. Because the p-n junction is spatially separated from the sawededge, no post-sawing processing is required.

Defects pose a serious problem to light emitting devices; they canreduce the optical efficiency as well as shorten the lifetime of thedevice. Often optical devices have dark line defects which multiplyduring operation. With continued operation, the density of these defectsincrease until the light output is reduced tt an unacceptable level anddevice failure occurs. In structures according to the present invention,wherein a small lattice mismatch exists between the substrate and theepitaxial layers , and a smaller mismatch between the GaN and AlGaNlayers, misfit dislocations due to these lattice mismatches are ofpotential concern.

These lattice mismatches may be addressed in three ways. The first isthe use of buffer layers including single composition layers as well assuperlattices and strain layer superlattices. The second method is toperform the growth on reduced area mesas which results in a reduction indislocation density in the epitaxial layers. Finally the choice of a SiCsubstrate, because of its high thermal conductivity, will keep thejunction temperature lower and result in much lower rate of defectpropagation. In addition, the closer thermal coefficient of expansionmatch between GaN and SiC compared to GaN and sapphire greatly reducesthe density of dislocations generated during cooldown after growth ofthe epitaxial layers.

As discussed above. the main lattice mismatch comes between thesubstrate and epitaxial layers. By choosing a SiC substrate instead ofthe more typical sapphire, the mismatch can be reduced from 13.8 toabout 0.9-3.4%, depending on the alloy composition. This will result ina large reduction in misfit dislocation density at thesubstrate/epitaxial layer interface. The misfit can be minimized byusing high Al alloy compositions in the cladding layer, which make upmost of the device volume. In addition, the closer TCE values willresult in fewer defects generated during cool-down after epitaxialgrowth. The TCE value for GaN is about 5.6 versus 5.1-5.9 for SiC (rangeof literature values) and about 7.5 for sapphire. Thus, the use of SiCinstead of sapphire will result in improved surface morphology, lesswafer warpage as well as a reduction in dislocation density due tomismatches in TCE. Still, the thicknesses required for these opticaldevices (LED, laser) will exceed the critical thickness for misfitdislocation generation and additional approaches are used to mitigatethese.

A buffer layer may be provided between the substrate and epitaxial layeras a manner of dealing with misfit dislocations in the growth of latticemismatched materials. Both AlN and GaN buffer layers have been used toadvantage with the nitrides on sapphire substrates. In addition, moresophisticated buffers, for example strained-layer (Al,Ga)Nsuperlattices, may be employed to further reduce the defect density.This type of buffer can be used in the III-V arsenides to reducedislocation densities by about a factor of 10.

A major problem with the growth of GaN on SiC substrates is cracking ofthe epi-layer. This is likely due to differences in the latticeconstants and thermal expansion coefficients of GaN and SiC, which causecracking when the samples are cooled from the growth temperature (˜1000°C.) to room temperature. Cracking of the GaN epi-layers is a severeproblem for the fabrication of electronic and optical devices from thesematerials.

As shown in Table III below (see "Properties of Group III Nitrides,"Edgar, J. H., INSPEC, London, 1994), GaN layers grown on sapphire arecompressed since the thermal coefficient of expansion (TCE) of GaN issmaller than that of sapphire. Conversely, GaN has a larger TCE than SiCso the layers are under tensile strain. The TCE of AlN is only slightlysmaller than that of SiC.

                  TABLE III    ______________________________________                                        Thermal                                        Expansion                       Lattice   Mismatch                                        Coefficient    Material  Symmetry Constant (Å)                                 to GaN (×10.sup.-6 /K)    ______________________________________    Wurtzite GaN              Hexagonal                       a = 3.189 --     5.59                       c = 5.185        7.75    Wurtzite AlN              Hexagonal                       a = 3.112 0.024  4.15                       c = 4.982        5.2    Basalplane              Hexagonal                       a = 4.758 0.138  7.5    Sapphire            c = 12.991      8.5    6H-SiC    Hexagonal                       a = 3.08  0.034  4.2                       c = 15.12        4.68    ______________________________________

SiC has a much closer lattice match to GaN (3.4%) than does sapphire.While an AlN layer grown at low temperature has been used as a bufferfor growth of GaN on SiC, improved crystal properties have been obtainedby growing a thin, approximately 1000 Å AlN buffer at temperaturesgreater than 1000° C. (T. W. Weeks, M. D. Bremser, K. S. Ailey, E.Carlson, W. G. Perry and R. F. Davis, GaN Thin Films Deposited viaOrganometallic Vapor Phase Epitaxy on 6H-SiC (0001) usingHigh-Temperature Monocrystalline AlN Buffer Layers, Appl. Phys. Lett. 67(1995) 401.)

It is noted that the mismatch between GaN and A1₂ O₃ in Table III is13.8%, not 33% which would be expected from the lattice constants,because the GaN rotates perpendicular to the c-axis to align 2110! GaNwith 0010! sapphire.

In the context of prior usage of thin AlN layers as buffers for GaNgrowth on SiC, it is noted that AlN has a closer lattice match andthermal expansion coefficient (TCE) to 6H-SiC than GaN. In actual growthof GaN layers on 10000 Å AlN buffer layers on SiC, the GaN layers crackdue to the larger differences in the TCE between GaN and AlN or SiC.

Thus, the parallel cracks that extend across the entire GaN layer, whenGaN is grown on SiC substrates using AlN buffers, are due to TCEdifferences between GaN and AlN (2.4%) and AlN and SiC (1.0%).

A transition buffer layer has been used to decrease the extent of GaNfilm cracking when growing on SiC (J. A. Edmond, V. Dmitriev, K. Irvine,Buffer Structure Between Silicon Carbide and Gallium Nitride andResulting Semiconductor Devices", U.S. Pat. No. 5,393,993). Thistransition buffer contains 3 layers which are described as a top(Al,Ga)N layer (%Al>%Ga) an intermediate (Al,Ga)N layer (%Ga>%Al) and abottom AlN layer. There is still, however, a large lattice mismatchbetween the GaN layer and the top (Al,Ga)N layer of the buffer (>1.4%)in this structure which introduces strain into the GaN layer. Among thethin AlN layers grown at high temperature that were used as buffers forGaN growth on SiC, according to this U.S. patent, was the buffer layerAl₀.90 Ga₀.10 N/Al₀.30 Ga₀.70 N/AlN.

The inventors herein have discovered a buffer structure which eliminatescracking, comprising a compositionally graded (Al,Ga)N buffer layerbetween the SiC substrate and the GaN epi-layer. More specifically, wehave found that improved GaN structural properties, as determined byx-ray double crystal rocking curve measurements, can be obtained byusing a high temperature AlN buffer. The full-width-at-half-maximum(FWHM) for the (0004) GaN reflection is 270 arc-sec with an AlN buffergrown at 1100° C. on 6H-SiC, compared to 428 arc-secs when AlN grown at550° C. is used as a buffer on 6H-SiC.

The inventors herein have demonstrated that crack-free GaN epi-layersseveral microns thick can be grown on 6H-SiC using an Al_(x) Ga_(1-x) Nbuffer layer where the Al composition (x) is graded from 1 at thebuffer-SiC interface to 0 at the GaN-buffer interface. The purpose ofthe compositional grading is to gradually vary the lattice constant andthermal expansion coefficient from that of AlN to that of GaN.

By contrast, a corresponding GaN layer grown on a 1000 Å thick AlNbuffer on 6H-SiC exhibited extensive film cracking.

It is to be noted here that the method of the present invention permitsthe low temperature step characteristic of prior art processes to beeliminated, so that growth temperatures on the order e.g., of1000°-1100° C. can be used for processing, As a result, the process ofthe present invention affords distinct advantages in terms of ease ofgrowth over the prior art, without the occurrence of cracking of theepi-layer.

FIG. 8 shows a GaN/SiC composite 70 including a GaN layer 71 grown on acompositionally graded (Al, Ga)N buffer layer 73 on a SiC substrate 72,where in the compositionally graded buffer, Al_(x) Ga_(1-x) N, x' variesfrom 1at the bottom to 0 at the top. The thickness of such graded buffercan range from 200 Å up to 5 μm, with a preferred range being from 500 Åup to 1 μm.

FIG. 9 shows a variation of the structure shown in FIG. 8, in which theGaN/SiC composite 78 includes a thin AlN buffer layer 76 is initiallygrown on SiC substrate 77 followed by the compositionally graded (Al,Ga)N buffer layer 75 and the GaN epi-layer 74. The thickness of the AlNbuffer layer 76 can range from 50 Å up to 5 μm, with a preferred rangebeing from 100 Å up to 1 μm.

In a structure of the type as shown in FIG. 8, with a 6 μm thick GaNepi-layer grown on a compositionally graded (Al, Ga)N buffer on 6H-SiC,no cracks were observed in the epi-layer. Additionally, the GaNepi-layers grown on compositionally graded (Al, Ga)N buffers on SiCexhibit superior structural and optical properties.

The buffer layers described above provide good morphological results,but because of the use of AlN, which is difficult to make conductive,the buffer is not conductive enough to use in devices where thesubstrate acts as one contact of the device, i.e., where current mustflow through the buffer layer.

According to another aspect of the invention, a conductive buffer layercan be fabricated as a variation of the compositionally graded bufferlayer described hereinabove. In such conductive buffer layer, theinitial Al_(x) Ga_(1-x) N layer starts with a lower Al composition layerthat can be doped to make it highly conductive in character, with x forthis starting composition being in the range of from x=1 to x=0.3,preferably with x between 0.7 to 0.4. The corresponding end compositioncould be GaN or a lower Al_(y) Ga_(1-y) N composition, where y<x. Theuse of the composition Al_(y) Ga_(1-y) N instead of GaN as the finalcomposition in the grade may be particularly useful when the epitaxiallayer to be grown on the buffer is (Al, Ga)N instead of GaN.

The conductive buffer layer described above may be made conductive bydoping with any suitable dopant species, so that the layer is of thedesired n-type or p-type. Illustrative of suitable n- and p-type dopantspecies which may be widely varied and usefully employed in the broadpractice of the present invention are silicon and magnesiun,respectively.

The buffer layers described hereinabove are made of (Al, Ga)Ncompositionally graded layers. They could also be made of (Al, In)N or(Al, Ga, In)N compositionally graded layers.

If the epitaxial layer to be grown on the buffer contains In, then itwill be advantageous to continue the compositionally graded buffer layerto the (In, Ga)N, (In. Ga)N, or (Al, Ga, In)N composition of theepitaxial layer. For example, if the epitaxial layer is In₀.1 Ga₀.9 N,then the buffer structure may illustratively begin with AlN and gradethrough (Al, Ga)N to GaN and then continue grading from GaN to In₀.1Ga₀.9 N.

A second defect reduction technique is to grow on reduced-area mesas.This technique has been used in the GaAs/InGaAs and Si/SiGe materialsystems. For example, a reduction in linear interface dislocationdensity with decreasing mesa diameter is observed for 7000 Å of In₀.05Ga₀.95 As grown on GaAs. A significant reduction occurs as the diameteris reduced from 200 to less than 50 μm. The mechanism for thisdislocation reduction is that, because of the small area, dislocationsare able to move to the edges of the pedestal and annihilate therebefore encountering and interacting with other dislocations. LEDstypically have an active region of about 170 μm in diameter and so weexpect some dislocation reduction by growing on a mesa of this diameter.

Lattice mismatch between a SiC substrate and the Ga*N epi-layer grown ontop of it introduces misfit dislocations into the epi-layer. Thesedefects degrade the structural, optical and electrical properties of theepi-layer and therefore have an adverse effect on the operatingcharacteristics of devices fabricated from the material To reduce theformation of defects in Ga*N epi-layers grown on SiC, and therebyimprove the structural and optical properties of Ga*N epi-layers grownon SiC substrates, reduced-area mesas are formed on the SiC substrateprior to epi-layers growth.

In Ga*N systems, small-area mesas reduce linear interface dislocationdensity. Due to the small area of the mesa, dislocations are able tomove to the edges of the pedestal and annihilate there beforeencountering and interacting with other dislocations. As a result, thestructural and optical properties of Ga*N epi-layers and quantum wellsgrown on reduced-area mesas on SiC substrates are significantlyimproved.

The present inventors have grown both GaN single layers and Al₀.15Ga⁰.85 N/GaN quantum well structures with GaN well widths of 150 Å and750 Å on reduced-area mesa SiC substrates. Growth on the reduced-areamesas results in an improved surface morphology of the GaN films. Thematerial on top of the mesa is much smoother than that of thesurrounding region.

FIG. 10 shows an individual mesa prepared as above. The mesas aretypically squares, with a side dimension ranging from 150 to 200 μm,with 160 to 180 microns preferred. Their "equivalent diameter" dimensiond ranges from 150 to 300 μm, with 250 to 300 μm preferred. As usedherein, the term "equivalent diameter" means the diameter of acylindrical shaped mesa which is equivalent in mesa top surface area tothe mesa being considered. Thus, when the mesa in fact is of cylindricalshape and circular cross-section, the equivalent diameter is equal tothe actual diameter measured on the circular top surface. For apolygonal or irregularly shaped mesa top surface, the top surface areamay be used to determine the diameter of a circular cross-section of thesame quantitative top surface area. The mesa height h relative to thesubstrate plane ranges from 1 to 15 μm, with 3 to 8 μm preferred.

FIGS. 11, 12, and 13 show a schematic representation of an exemplaryprocess used to produce such mesa-defined devices. The substrate 61 isfirst patterned with mesas 62 of the desired shape and dimension. Mesasmay preferably be either circular or square, with circular preferred,having 150-300 μm equivalent diameter, preferably in the range of200-250 μm. The substrate is patterned with metal (Ni, Al) which hasbeen photolithographically patterned to the shape of the desired mesas.The non-masked areas are etched using reactive ion etching to a depth of1-15 μm, and preferably 3-8 μm. The metal is then removed and thesubstrate is cleaned for epitaxial growth. The LED structure, whichconsists of layers 63, the active layers and cladding, is formed bysuccessive steps of growth (described above). Again, the structures aremasked with metal to protect the mesa tops, and to remove the materialwhich was conformally deposited on the mesa sidewalls, another reactiveion etching step is performed. The metal may form the top contact 64 orit can constitute the mask, in which case another step is required toform the contact. A back metal contact 65 is deposited on the back ofthe wafer. After separation by a suitable technique, e.g. dicing,mechanical or laser scribing and cleaving, or sawing, the individualdevices 66 are ready to be packaged.

The third approach to reducing the impact of defects arises from the useof the SiC substrate. Dark line defects are known to multiply because ofdevice heating during operation. One way to reduce this effect is byreducing heating of the device. The SiC substrate has a thermalconductivity that is about 10 times greater than that of sapphire.Therefore, the junction region will be cooler, which will reduce thepropensity of the defects to multiply, and thus extend device lifetime.

Methods for forming the device structures can be adapted from techniquesthat have been developed in SiC, GaAs and silicon materials and deviceprocessing technology. Single crystal silicon carbide substrates can beprepared by sublimation growth, as described by Tairov and Ziegler (Y.M. Tairov, V. F. Tsvetkov, J. Crystal Growth 43 (1978) 209; 52 (1981)146; G. Ziegler et al., IEEE Trans. Electron Devices ED-30 (1983) 277)).GaN and other III-V nitride layers may be deposited by severaltechniques, including low and high pressure and plasma-enhanced chemicalvapor deposition (CVD), reactive ionized-cluster beam deposition,reactive and ionized molecular beam epitaxy (MBE), and hydride vaporphase epitaxy (VPE). For n-type materials, doping of the III-V nitrideis achieved typically using silicon from silane or disilane. Othern-type dopants include sulfur, selenium and tellurium. P-type doping hasbeen more problematic in the III-V nitrides. One recent method is to useelectron beam irradiation to activate dopants. The disadvantage of thisapproach is that it is a slow, expensive process and cannot be used instructures with arbitrary thickness.

The fabrication method of the present invention employs a combination ofcontrol of specific growth conditions as well as thermal annealing aftergrowth to achieve active p-type dopants. Typical p-type dopants includeMg, Be, Zn, Cd and C.

FIG. 14 shows a schematic representation of a semiconductor laser 90according to the present invention. The laser comprises a top contact91, top Bragg reflector 92, bottom Bragg reflector 93, cladding layer94, active layer 95, cladding layer 96, buffer 97, substrate 98, bottomcontact 99, and passivation layer 100. Other semiconductor laserstructures can incorporate a key aspect of the invention, the green-blueto ultraviolet light emitting Ga*N material on a base structurecomprising a SiC substrate selected from the group consisting of 2H-SiC,4H-SiC and a-axis oriented 6H-SiC. Layers 94 and 96 have larger bandgaps than layer 95, and are complementarily doped, one p-type and onen-type.

FIG. 15 depicts another example of a laser structure 110 according tothe present invention. Laser 110 comprises top contact 111, contactlayer 112, cladding layer 113, active layer 114, cladding layer 115,substrate 116, and bottom contact 117. Face 118 and the correspondingface on the opposite side of the device constitute laser facets ormirrors. Light is emitted from the active layer on the facet sides,beneath the contact. Layers 113 and 115 have larger band gaps than 114,and are complementarily doped, one p-type and one n-type. Although notshown in the specific embodiment illustrated in FIG. 15, it is to beappreciated that the laser structure shown in FIG. 15 could employ abuffer layer between the cladding layer 115 and the substrate 116.

As discussed hereinabove, the present invention contemplates the use ofbuffer layers in Ga*N structures such as Ga*N/SiC semiconductor devicesor precursors thereof, in which the buffer layers function to providethe overall structure with a good lattice match and a good coefficientof thermal expansion match between the layer of single crystal siliconcarbide and the layer of gallium nitride.

The invention thus realates to gallium nitride semiconductor devices orprecursor articles therefor, in which the device or precursor comprisesa layer of single crystal silicon carbide and a layer of single crystalgallium nitride, having a buffer layer therebetween comprising acompositionally graded Ga*N layer.

Such compositionally graded Ga*N layer may suitably comprise acompositionally graded Al_(x) Ga_(1-x) N buffer layer between thegallium nitride and silicon carbide layers, wher e in x can range from 0to 1, and wherein the buffer layer is compositionally graded from aninterface of the the buffer layer with the silicon carbide layer atwhich x is 0, to an interface of the buffer layer with the galliumnitride layer at which x is 1. The buffer layer thus is varied acrossits thickness between the GaN and SiC layers sandwiching the bufferlayer therebetween, so that the buffer layer at its respectiveextremeties provides good epitaxial and TCE match to the contiguous GaNand SiC layers bounding the buffer material.

Preferably, such buffer layer is graded smoothly from an initial valueof x at the interface of the buffer layer with the SiC layer, to a finalvalue of x at the other interface of the buffer layer with the GaNlayer, with continuous progression of the stoichiometry of the bufferlayer metal constiuents along the thickness thereof, however it is alsopossible in the practice of the invention to utilize a buffer layerwhich is graded in stepwise or incremential fashion along its thickness,or the variation of the stoichiometric change may be intermediate, or acombination of, such continous and step changes.

Thus, the invention contemplates a semiconductor device or precursorthereof, comprising a silicon carbide substrate and an epitaxial layerof gallium nitride, having a buffer layer therebetween comprising acompositionally graded (Al,Ga)N buffer layer. The SiC substrate maycomprise 6H-SiC or other SiC polytypes, such as 2H-SiC or 4H-SiC, withthe Al_(x) Ga_(1-x) N buffer layer wherein the Al composition (x) isgraded from 1 at the buffer-SiC interface to 0 at the GaN-bufferinterface. The thickness of such graded buffer layer may suitably rangefrom 200 Å up to 5 μm, with a preferred range being from 500 Å up to 1μm.

A variation of the above-described semiconductor device or precursorstructure may comprise a thin AlN buffer layer which is initially grownon SiC followed bv the compositionally graded (Al, Ga)N buffer and theGaN epitaxial layer. The thickness of such AlN buffer layer can rangefrom 50 Å up to 5 μm, with a preferred range being from 100 Å up to 1μm.

As a variation of the above-described semiconductor device and precursorstructures, the conductive buffer layer of the formula Al_(x) Ga_(1-x) Nmay include a lower Al composition layer that is doped to render ithighly conductive in character, such that x is in the range of from x=1to x=0.3, preferably with x being be tween 0.7 to 0.4. The correspondingend composition could be GaN or a lower Al_(y) Ga_(1-y) N composition,wherein y<x. The use of the composition Al_(y) Ga_(1-y) N instead of GaNas the final composition in the graded buffer layer may be particularlyuseful when the epitaxial layer to be grown on the buffer is (Al, Ga)Ninstead of GaN.

Any suitable dopant species and dopant concentrations may be employedwithin the skill of the art, to provide the buffer layer with a desiredconductive character and functional operating characteristics. Forexample, the buffer layer may be doped to be n-type, by use of siliconas a dopant species. As another example, the buffer layer may be dopedto be p-type, by use of magnesium as a dopant species.

More generally, the Ga*N buffer layers broadly described hereinabove maybe formed of any suitable buffer layer material, such as an (Al, Ga)Ncompositionally graded layer, an (Al, In)N compositionally graded layer,or an (Al, Ga, In)N compositionally graded layer.

As yet another example, if an epitaxial layer to be grown on the buffercontains In, the compositionally graded buffer layer may be continued tothe (In, Ga)N, (In, Ga)N, or (Al, Ga, In)N composition of the epitaxiallayer, e.g., if the epitaxial layer is In₀.1 Ga₀.9 N, then the bufferstructure may illustratively begin with AlN and grade through (Al, Ga)Nto GaN, and then continue grading from GaN to In₀.1 Ga₀.9 N.

FIG. 16 shows a schematic sectional elevation view of a light emittingdiode 50 fabricated with a gallium nitride active layer, grown on acompositionally graded (Al, Ga)N buffer layer 54, on a silicon carbidesubstrate. The LED 50 comprises a green-blue to ultraviolet lightemitting GaN material 51 on a base structure 53 comprising a SiCsubstrate. The diode structure includes a p-n junction comprising GaNlayers 51 and 52, a contact 55 on the upper surface of the GaN layer 52,and a contact 56 on the bottom surface of the SiC substrate basestructure 53.

It will be appreciated that other of the Ga*N devices and structuresdescribed herein could also advantageously utilize a graded buffer layerintermediate the Ga*N layer and the SiC layer thereof.

The features and advantages of the invention are more fully shown by thefollowing non-limiting, illustrative examples.

EXAMPLE I

SiC substrates were prepared with reduced-area mesas for subsequent GaNgrowth. The substrates were (0001) 6H-SiC with a misorientation 4°towards <1120>.

The fabrication process included the following process steps:

1. Lithography to define mesas on the SiC substrate..

2. Evaporation and lift-off of Ni to define circular mask area foretching.

3. Reactive ion etching of SiC to define mesas.

4. Removal of Ni and cleaning of SiC substrate.

Scanning electron microscopy (SEM) of reduced-area mesas etched in theSiC substrate showed the mesa height to be ˜3 μm and the mesa diametersto vary from about 25 to 300 μm.

EXAMPLE II

GaN single layers and Al₀.15 Ga₀.85 N/GaN quantum well structures withGaN well widths of 150 Å and 750 Å were grown on reduced-area mesa SiCsubstrates. Growth on the reduced-area mesas resulted in an improvedsurface morphology of the GaN films. The material on top of the mesa wasmuch smoother than that of the surrounding region.

EXAMPLE III

AlGaN/GaN quantum well structures were grown on reduced area mesa SiCsubstrates. Structure A consisted of a 750 Å thick GaN active regionsandwiched between a 2000 Å Al₀.15 Ga₀.85 N cap and a 3 μm thick Al₀.15Ga₀.85 N cladding layer. Structure B consisted of a 150 Å thick GaNactive region between two 300 Å thick Al₀.04 Ga₀.96 N layers between a2000 Å Al₀.15 Ga₀.85 N cap and a 3 μm thick Al₀.15 Ga₀.85 N claddinglayer.

Photoluminescence (PL) measurements were used to characterize theoptical properties of the AlGaN/GaN hetero structures. Thesemeasurements were carried out at room temperature using a 325 nm He-Cdlaser as the excitation source. Focusing optics were used to obtain alaser spot size on the sample of approximately 100 μm in diameter.

The GaN band-edge PL intensity from material grown on top of the mesaswas about a factor of two greater than from off-mesa regions of the 3 μmthick GaN layer and the 750 Å GaN well sample, as shown in FIGS. 17 and18, respectively. The GaN band-edge peak is located at 364 nm in the GaNlayer (FIG. 17) but is shifted ˜28 meV higher in energy to 361 nm forthe 750 Å thick GaN well (FIG. 18). The 750 Å GaN well is too thick toattribute the shift to quantum confinement effects. The shift in energyis likely due instead to strain resulting from the 0.4% lattice mismatchbetween the GaN and Al₀.15 Ga₀.85 N layers.

The effect of the reduced area mesa was most pronouned for the 150 Å GaNquantum well sample. As shown in FIG. 19, GaN band-edge PL was onlyobserved from material grown directly on top of the mesas. The GaNband-edge peak from the 150 Å GaN quantum well is at 359 nm, shifted ˜47meV higher in energy than for the thick GaN layer. The PL intensity near340 nm in FIGS. 18 and 19 is from the AlGaN cladding layers.

The increased PL intensity from GaN and AlGaN/GaN quantum wells islikely due to an improvement in the structural properties of thematerial grown on the reduced area mesas. The surface morphology of themesas was, in general, less rough than that of the surrounding regions,consistent with a decrease in the defect density of the material. Thesmoother surface may in part be responsible for the increased intensityon the mesas.

EXAMPLE IV

Attempts were made to grow GaN on the following buffers grown at hightemperature: 1000 Å AlN, 1000 Å AlGaN and 1000 Å AlN/1000 Å Al₀.10Ga₀.90 N/1000 Å AlN. All of these gave good x-ray diffraction (XRD) FWHMvalues in the 200 arcsec range, but the GaN epi-layers were cracked.

EXAMPLE V

A buffer structure which was determined to eliminate cracking of the GaNepi-layers comprises a compositionally graded (Al,Ga)N buffer whereinthe Al content is graded from AlN at the substrate to GaN at the top.The growth rate of AlN is much lower than that of GaN for a givenmetalorganic (MO) mole fraction.

The initial intent of the work in this example was to try and keep thegrowth rate constant throughout the buffer layer, and thus an initialTMAl flow rate and a final TMGa flow rate were chosen that would givethe same growth rate and then would permit grading the flows to zero.This is shown in FIG. 20 wherein the flow rates are plotted as afunction of time to show the variation of TMAl and TMGa flows duringgrowth of a buffer.. The total MO flow rate (TMAl+TMGa) decreased duringthe growth of the buffer. The growth rate and composition gradient ofthe buffer is not known. The growth period was chosen to yield athickness of ˜1000-2000 Å, based on the initial AlN and final GaN growthrates. Cracking of the GaN film was inhibited.

GaN growth on the graded (Al,Ga)N buffer was repeated four times. In allinstances, all of the GaN layers were free of cracks. The XRD FWHM ofthe GaN peak varied depending of the FWHM of the SiC substrate peak, asshown in Table IV below. The room temperature PL peak width did notsignificantly change.

                  TABLE IV    ______________________________________    Run No. GaN XRD FWHM SiC XRD FWHM                                     PL peak width    ______________________________________    A       110 arcsec    42 arcsec  38Å    B       147          129         38Å    C       281          208         38Å    D       181          135         36Å    ______________________________________

EXAMPLE VI

A GaN layer was grown on a compositionally graded (Al, Ga)N buffer on6H-SiC. A double crystal x-ray rocking curve of this sample was made,and is shown in FIG. 21. The full width at half-maximum (FWHM) of theGaN epi-layer is 110 arcsecs, indicating very good crystal quality. TheSiC FWHM is 42 arcsecs. The room temperature photoluminescence of thissample contained a GaN band-edge peak at 363 nm with a peak of 36 nmwith a peak width of 36 meV.

EXAMPLE VII

Reflectivity Measurements of III-V Nitride Bragg Reflector Stacks

GaN/Al₀.15 Ga₀.85 N and Al₀.12 Ga₀.88 N/Al₀.40Ga₀.60 N Bragg reflectorstacks were fabricated on both SiC and sapphire substrates. FIG. 22compares the simulated and experimentally measured reflectivity from a24 repeat period 42.8 nm GaN/44.6 nm Al₀.15 Ga₀.85 N Bragg reflector.The predicted reflectivity is 80% at 430 nm compared to theexperimentally measured reflectivity of 80% at 438 nm.

EXAMPLE VII

Photolumninescence Measurements of Bragg Reflectors

Photoluminescence measurements were used to study and compare theoptical properties of Bragg reflector stacks grown on SiC and sapphiresubstrates. The spectra were obtained at room temperature using focusedlight from a pulsed nitrogen laser emitting at 337.1 nm. PL spectra ofthe Bragg reflectors grown on SiC and sapphire substrates are shown inFIGS. 23 and 24, respectively. FIG. 23 is a graph of PL spectra as afunction of input pump power for a GaN/AlGaN Bragg reflector stack grownon a SiC substrate. FIG. 24 is graph of PL spectra as a function ofinput pump power for a GaN/AlGaN Bragg reflector stack grown on asapphire substrate.

Spectra obtained from the Bragg reflector grown on SiC for low powersexhibit a single feature peaked at 363 nm. At higher powers, a sharppeak, corresponding to stimulated emission, at 372 nm emerges from thebroader band-edge peak and increases rapidly with pump intensity. Thepeak intensity at 372 nm clearly exhibits a threshold behavior with anapproximate threshold pump density of 70 MW/cm². The threshold pumpdensity for the onset of stimulated emission is much higher in the Braggreflector stack grown on sapphire, as shown in FIG. 24. In this case.the threshold pump density is 200 MW/cm².

EXAMPLE VIII

VCSEL Structures utilizing Bragg Reflectors

III-V nitride Bragg reflectors were employed in the fabrication ofvertical cavity surface emitting laser (VCSEL) structures. The VCSELstructure, shown schematically in FIG. 25, consisted of a 10 μm GaNactive region within a vertical cavity formed by 30-period Al₀.40 Ga₀.60N-Al₀.12 Ga₀.88 N Bragg reflector (397 Å/372 Å) multilayer stacks. TheBragg reflectors were designed to have 99% reflectivity at 370 nm. TheIII-V nitride layers were grown by low pressure MOVPE in separate growthruns on a c-plane (0001) sapphire substrate and an (0001) 6H-SiCsubstrate miscut 4° toward the <1120>. Trimethylgallium,trimethylalurninurn, and ammonia were used as precursors with hydrogenas the carrier gas. Optimized buffers were initially grown on each typeof substrate followed by the growth of an Al₀.12 Ga₀.88 N layer and theVCSEL epilayers at 1100° C.

Optical pumping experiments were performed using a pulsed nitrogen laseras the pump source. The VCSELs were photopumped from the side, with thepump beam focused to a ˜70 μm spot on an edge of the sample, and surfaceluminescence was collected by a lens system and focused into an 0.64 mmonochromator with a UV grating. Peak intensities of the focused pumplight were estimated from the spot size (70 μm), the pulse width, andthe pulse energy. The pulse width is taken to be 0.8 ns, as given by themanufacturer specifications. Data were determined for pump intensitiesof 0.65-4.3 MW/cm², corresponding to measured pulse energies of 0.94-6.3μJ/pulse.

FIG. 26 shows double crystal x-ray rocking curves (w-scan) of VCSELgrown on a) 6H-SiC and b) sapphire substrates. The structural quality ofthe VCSEL layers was characterized using double crystal x-raydiffraction. The full-width-at-half maximum (FWHM) for the (0004) GaNreflection in FIG. 26, in curve (a) is 260 arcsecs, while the zero-orderreflection from the top and bottom Al₀.40 Ga₀.60 N-Al₀.12 Ga₀.88 Nmultilayer stacks had a FWHM of 580 arcsecs. It was not possible toclearly resolve higher-order multilayer reflections. likely due to thebroadness of the zero-order peak. Similar peak positions and widths weremeasured for the VCSEL grown on sapphire, as shown in curve (b) in FIG.26.

Surface emission spectra obtained at three different pump intensitiesare shown in FIG. 6 for the VCSEL grown on sapphire. At a pump intensityof 1.3 MW/cm² (˜0.65 P_(th)), the spectrum consists of a single PL peakcentered at 368 nm. At an input intensity of 2.2 MW/cm² (1.1 P_(th)),the emergence of a sharp feature peaked near 363.5 nm is evident. Theintensity of this peak becomes dominant at higher powers and isaccompanied by the emergence of additional peaks at 362.1 nm and 364.8nm, which are equally separated from the dominant mode by ˜1.3 nm. Peakand integrated intensities of the emission spectra are plotted in FIG. 7as a function of pump intensity. The peak intensity increases slowlywith pump intensity up to ˜2 MW/cm², above which the intensity increasesmuch more rapidly as the sharp 363.5 nm line emerges and dominates thespontaneous emission peak. Estimation of lasing threshold fromextrapolation of both the peak and integrated intensity data yieldsP_(th) ˜2.0 MW/cm².

The observed mode spacing of ˜1.3 nm is consistent with expectations forthe VCSEL structure. The structure was modeled using a transfer matrixapproach which allows for dispersion and loss in the layer materials.Wavelength-dependent refractive indices for the various layers wereobtained from parameter interpolation and an empirical model whichclosely fits near-gap ellipsometry data for GaN and Al₀.1 Ga₀.90 N. Themodel predicts a mode spacing of 1.1 nm for the VCSEL and a peakreflectivity near the design wavelength of 370 nm for the Braggreflectors.

FIG. 27 is a graph of room temperature surface emission spectra forVCSEL grown on sapphire at increasing pump intensities. FIG. 28 is agraph of peak and integrated intensities of the VCSEL/sapphire emissionspectrum as a function of pump intensity. The threshold pump intensity,extrapolated from data on peak emission intensities, is 2.0 MW/cm².

FIG. 29 is a graph of the spectral characteristics of surface emissionfrom the VCSEL grown on SiC both at (P_(th)) and above the threshold(1.7 P_(th)), showing the room temperature emission spectra for VCSELgrown on SiC at increasing pump intensities. In this case, a singlenarrow (0.2 nm) emission peak at 367.5 nm was observed, in contrast tothe multi-mode spectrum obtained from the VCSEL grown on sapphire. Theintensity of the 367.5 nm peak increases rapidly with pump power above athreshold value of ˜3 MW/cm². While several areas along the length ofeach sample were photopumped, single mode lasing was only observed fromthe VCSEL grown on the SiC substrate. The emission from the VCSEL grownon SiC is at a higher wavelength (367.5 nm) than the VCSEL grown onsapphire (363.5 nm), consistent with the shift in GaN near band-edgespontaneous emission that we have previously observed for GaN grown onthe two types of substrates.

PREFERRED MODES OF CARRYING OUT THE INVENTION

The invention is advantageously carried out in LED, laser and otherdevice applications with SiC substrates providing high charge carriermobility, to maximize device performance.

Accordingly, 4H-SiC or 2H-SiC may be employed as polytypes for the SiCsubstrate due to their high intrinsic mobility characteristics.Alternatively, where the SiC substrate has a more anisotropic characterand a lower intrinsic charge carrier mobility, such as is the case withthe 6H-SiC polytype, the substrate is desirably deployed with thesubstrate crystal's highest mobility axis (a-axis in the case of 6H-SiC)being aligned with the current flow direction in the device.

In applications in which broad-spectrum transparency is desired forlight emission, larger band gap materials such as 2H-SiC and 4H-SiCpolytypes afford significant optical advantages. For high-energytransmission applications such as communications and spectroscopy, the2H-SiC and 4H-SiC polytypes are highly desirable, in consequence oftheir large band gap and low cut-off wavelength characteristics.

Devices according to the invention having quantum well active regionsmay also fabricated in which the well width is selectively dimensionedand the well composition and barrier layer composition are selectivelyprovided, to yield a specific emission wavelength and energy.

Devices fabricated in accordance with the invention may employ adielectric Bragg mirror comprising a series of metallonitride materiallayers under the light emitting structure, to increase the lightemission of the device. The inherent series resistances and chargecarrrier potential barriers of Bragg mirror structures which degradetheir optical and performance properties can be ameliorated by theprovision of buffer layers (an intermediate composition, short periodsuperlattices, or strained layer superlattices) between the successivemain layers of the structure.

The fabrication of light-emitting optical devices in the practice of theinvention may be advantageously conducted without having to carry outetching after dicing has been performed on the device structure, by themethod of etching mesas on the SiC substrate prior to growth of Ga*Nepilayers, with passivation of the mesa edge, prior to formation ofcontacts, so that the active region of the device is removed from thedicing region. Such spatial separation and isolation of the active areaof the device, e.g., the p-n junction of a LED structure thus preventsthe occurrence of dicing damage which may significantly degrade thelight emission efficiency or service life of the device. The passivationlayer may be formed of a material such as silicon dioxide or siliconnitride, deposited by CVD, sputtering, plasma-assisted deposition, orother suitable technique, at a thickness of for example 200 to 2000 Å.

A preferred fabrication technique usefully employed in the practice ofthe invention involves utilizing reduced area mesas for Ga*N devicefabrication. By such technique, the interfacial dislocation density issignificantly reduced, relative to other substrate base structures, as aresult of the freedom of dislocations to move to edge regions of thesmall sized mesas. where dislocations are annihilated.

In consequence, the devices fabricated on such small sized mesas havesubstantially improved structural and optical properties, includingincreased intensity of emitted light, and smooth surface morphology ofthe mesa areas.

The preferred mesa structures may include mesas with a height of 1-15 μmand an equivalent diameter of the mesa top surface area in the range offrom 25-300 μm, on a suitable SiC substrate, such a for example (0001)6H-SiC with a misorientation 4° towards <1120>. The specific process ofmesa formation may illustratively comprise the steps of (1) lithographyto define mesas, (2) evaporation and lift-off of Ni to define circularmask areas for etching, (3) reactive ion etching of SiC to define mesas,and (4) removal of Ni and cleaning of SiC substrate.

Epi-layer crack-resistant Ga*N structures such as Ga*N/SiC semiconductordevices or precursors thereof, may be fabricated in which the bufferlayers are employed to provide the overall structure with a good latticematch and a good coefficient of thermal expansion match silicon carbideand gallium nitride layers. Such buffer layer advantageouslycomprises acompositionally graded Ga*N layer, such as a compositionally gradedAl_(x) Ga_(1-x) N buffer layer between the gallium nitride and siliconcarbide layers, wherein x can range from 0 to 1, and wherein the bufferlayer is compositionall, and preferably smoothly, graded from aninterface of the the buffer layer with the silicon carbide layer atwhich x is 0, to an interface of the buffer layer with the galliumnitride layer at which x is 1. The thickness of such graded buffer layerranges from 200 Å up to 5 μm, preferably from 500 Å up to 1 μm.

Alternatively, the semiconductor device or precursor structure maycomprise a thin AlN buffer layer which is initially grown on SiCfollowed by the compositionally graded (Al, Ga)N buffer and the GaNepitaxial layer, wherein the thickness of such AlN buffer layer rangesfrom 50 Å up to 5 μm, preferably from 100 Å up to 1 μm.

The semiconductor device or precursor structure may alternativelycomprise a conductive buffer layer of the formula Al_(x) Ga_(1-x) Nincluding a lower Al composition layer that is doped with suitabledopant species (e.g., Si, Mg) to render it highly conductive incharacter, such that x is in the range of from x=1 to x=0.3, preferablybetween 0.7 to 0.4. The corresponding end composition could be GaN or alower Al_(y) Ga_(1-y) N composition, wherein y<x, and the use of thecomposition Al_(y) Ga_(1-y) N instead of GaN as the fmal composition inthe graded buffer layer may be particularly useful when the epitaxiallayer to be grown on the buffer is (Al, Ga)N instead of GaN.

INDUSTRiAL APPLICABILITY

The light emitting optical devices of the present invention have utilityin industrial devices such as lasers and LEDs, for applications such ashigh density optical storage, full color displays, color determinationsystems, and spectroscopic analysis sources. In such applications, theelectroluminescent devices of the invention efficiently emit very brightlight, as measured by the indicia of quantum efficiency and luminousintensity.

We claim:
 1. A semiconductor device or precursor thereof, comprising alayer of single crystal silicon carbide and a layer of single crystalgallium nitride, having a buffer layer therebetween comprising acompositionally graded Ga*N layer.
 2. A semiconductor device orprecursor thereof according to claim 1, wherein the silicon carbidecomprises a polytype selected from the group consisting of 2H-SiC,4H-SiC, and 6H-SiC.
 3. A semiconductor device or precursor thereofaccording to claim 1, wherein said compositionally graded Ga*N layercomprise a compositionally graded Al_(x) Ga_(1-x) N buffer layer betweenthe gallium nitride and silicon carbide layers, wherein x can range from0 to 1, and the buffer layer is compositionally graded from an interfaceof the the buffer layer with the silicon carbide layer at which x is 0,to an interface of the buffer layer with the gallium nitride layer atwhich x is
 1. 4. A semiconductor device or precursor thereof, comprisinga silicon carbide substrate and an epitaxial layer of gallium nitride,having a buffer layer therebetween comprising a compositionally graded(Al,Ga)N buffer layer.
 5. A semiconductor device or precursor thereofaccording to claim 4, wherein said buffer layer comprises an AlGal-,Nbuffer layer where the Al composition (x) is graded from 1 at thebuffer-SiC interface to 0 at the GaN-buffer interface.
 6. Asemiconductor device or precursor thereof according to claim 4, whereinthe thickness of said buffer layer is from 200 Å to 5 μm.
 7. Asemiconductor device or precursor thereof according to claim 4, whereinthe thickness of said buffer layer is from 500 Å to 1 μm.
 8. Asemiconductor device or precursor thereof, comprising a silicon carbidesubstrate and an epitaxial layer of gallium nitride, having a buffermaterial therebetween comprising an AlN buffer layer on the siliconcarbide substrate, and a compositionally graded (Al,Ga)N buffer layerbetween the AlN buffer layer and the gallium nitride epitaxial layer. 9.A semiconductor device or precursor thereof according to claim 8,wherein the thickness of said AlN buffer layer is from 50 Å to 5 μm. 10.A semiconductor device or precursor thereof according to claim 8,wherein the thickness of said AlN buffer layer is from 100 Å to 1 μm.11. A semiconductor device or precursor thereof, comprising a siliconcarbide substrate and an epitaxial layer of gallium nitride, having abuffer material therebetween comprising a conductive buffer layer of theformula Al_(x) Ga_(1-x) N wherein the buffer layer includes a lower Alcomposition layer that is doped to render it highly conductive incharacter.
 12. A semiconductor device or precursor thereof according toclaim 11, wherein x is in the range of from x=1 to x=0.3.
 13. Asemiconductor device or precursor thereof according to claim 11, whereinx is in the range of from x=0.7 to x=0.4.
 14. A semiconductor device orprecursor thereof according to claim 11, wherein the buffer materialcomprises an end composition adjacent said gallium nitride epitaxiallayer, selected from the group consisting of GaN and a lower Al_(y)Ga_(1-y) N composition, where y<x.
 15. A semiconductor device orprecursor thereof according to claim 14, wherein the buffer materialcomprises a s said end composition Al_(y) Ga_(1-y) N and the galliumnitride epitaxial layer comprises (Al, Ga)N.
 16. A semiconductor deviceor precursor thereof according to claim 11, wherein the conductivebuffer layer is an n-type material.
 17. A semiconductor device orprecursor thereof according to claim 11, wherein the conductive bufferlayer is a p-type material.
 18. A semiconductor device or precursorthereof, comprising a silicon carbide substrate and an epitaxial layerof gallium nitride, having a buffer material therebetween comprising abuffer layer selected from the group consisting of (Al, Ga)Ncompositionally graded layers, (Al, In)N compositionally graded layers,and (Al, Ga, In)N compositionally graded layers.
 19. A semiconductordevice or precursor thereof according to claim 18, wherein the epitaxiallayer of gallium nitride contains In.
 20. A semiconductor device orprecursor thereof according to claim 19, wherein the epitaxial layer ofgallium nitride comprises an epitaxial layer formed of material selectedfrom the group consisting of (In, Ga)N, (In, Ga)N, and (Al, Ga, In)N.21. A semiconductor device or precursor thereof according to claim 19,wherein the epitaxial layer of gallium nitride comprises In₀.1 Ga₀.9N,and the buffer material includes AlN grading through (Al, Ga)N to GaN,and grading from GaN to In₀.1 Ga₀.9 N.
 22. A semiconductor device orprecursor thereof, comprising a SiC substrate element including amesa-shaped portion having a Ga*N quantum well structure or other Ga*N structure formed thereon, where in t he mesa-shaped portion has anequivalent diameter in the range of from 50 to 300 μm, and a height offrom 1 to 15 μm.
 23. A semiconductor device or precursor thereofaccording to claim 22, wherein the SiC substrate comprise an SiCpolytype selected from the group consisting of 2H-SiC, 4H-SiC, anda-axis oriented 6H-SiC.
 24. A semiconductor device or precursor thereofaccording to claim 22, wherein said mesa-shaped portion has a quantumwell structure thereon, and said quantum well structure includes a wellhaving a well width of from 150 Å to 750 Å.
 25. A method for fabricatinga light emitting diode comprising the steps of:a. forming an M_(x)M'_(y) M"_(1-x-y) N barrier layer on a conductive substrate, wherein M,M', and M" are metals compatible with nitrogen (N) in the compositionM_(x) M'_(y) M"_(1-x-y) N, and the composition M_(x) M'_(y) M"_(1-x-y) Nis stable at room temperatures and pressures; b. forming an M_(a) M'_(b)M"_(1-a-b) N active region on the M_(x) M'_(y) M"_(1-x-y) N barrierlayer; c. forming an M_(c) M'_(d) M"_(1-c-d) N barrier layer on theM_(a) M'_(b) M"_(1-a-b) N active region; d. forming a heavily dopedcontact layer on the M_(c) M'_(d) M"_(1-c-d) N barrier layer; and e.forming a first contact on the heavily doped contact layer and a secondcontact on the substrate, wherein x, y, a, b, c, and d are numbers, eachnumber having a magnitude between zero and one and wherein x+y, a+b, andc+d, each have a magnitude between zero and one.
 26. The method of claim25 wherein the substrate is a selected polytype of silicon carbide (SiC)selected from a group of polytypes of SiC including 2H-SiC, 4H-SiC, anda-axis oriented 6H-SiC.
 27. The method of claim 26 wherein the SiCsubstrate is oriented to have a high mobility axis aligned with thedirection of electron current.
 28. A method for fabricating a lightemitting diode comprising the steps of:a. forming a first barrier layeron a conductive substrate, the first barrier layer being a firstselected composition; b. forming an active region of a second selectedcomposition on the first barrier layer; C. forming a second barrierlayer of a third selected composition on the active region; d. forming aheavily doped contact layer of a fourth selected composition on thesecond barrier layer; and e. forming a first contact on the heavilydoped contact layer and a second contact on the substrate, wherein thefirst, second, third and fourth selected compositions are selected froma group of compositions M_(x) M'_(y) M"_(1-x-y) N, M_(a) M'_(b)M"_(1-a-b) N, M_(c) M'_(d) M"_(1-c-d) N and M_(e) M'_(f) M"_(1-e-f) N,respectively, wherein x, y, a, b, c, d, e, and f are numbers, eachnumber having a magnitude between zero and one and wherein x+y, a+b,c+d, and e+f, each has a magnitude between zero and one.
 29. The methodof claim 28 wherein the substrate is a selected polytype of siliconcarbide (SC) selected from a group of polytypes of SC including 2H-SiC,4H-SiC, and a-axis oriented 6H-SiC.
 30. The method for fabricating alight emitting diode with a Bragg mirror, the method comprising thesteps of:a. forming a Bragg mirror including alternating layers of AlN,GaN, InN, or their alloys on a conductive substrate; b. forming an M_(x)M'_(y) M"_(1-x-y) N barrier layer on a conductive substrate, wherein M,M', and M" are metals compatible with nitrogen (N) in the compositionM_(x) M'_(y) M"_(1-x-y) N, and the composition M_(x) M'_(y) M"_(1-x-y) Nis stable at room temperatures and pressures; c. forming an M_(a) M'_(b)M"_(1-a-b) N active region on the M_(x) M'_(y) M"_(1-x-y) N barrierlayer; d. forming an M_(c) M'_(d) M"_(1-c-d) N barrier layer on theM_(a) M'_(b) M"_(1-a-b) N active region; e. forming a heavily dopedcontact layer on the M_(c) M'_(d) M"_(1-c-d) N barrier layer; and f.forming a first contact on the heavily doped contact layer and a secondcontact on the substrate, wherein x, y, a, b, c, and d are numbers, eachnumber having a magnitude between zero and one and wherein x+y, a+b, andc+d, each have a magnitude between zero and one.
 31. The method of claim30 wherein the substrate is a selected polytype of silicon carbide(SiC), selected from a group of polytypes of SiC including 2H-SiC,4H-SiC, and a-axis oriented 6H-SiC.
 32. The method for fabricating lightemitting diodes (LEDs) comprising the steps of:a. forming a firstbarrier layer on a conductive substrate; b. forming an active region onthe first barrier layer; C. forming a second barrier layer on the activeregion; d. forming a heavily doped contact layer on the second barrierlayer; e. masking each LED device area; f. mesa etching trenches betweenthe LED device areas and leaving the LEDs as mesas separated by thetrenches; g. passivating sidewalls of the LEDs; h. forming a firstcontact on the heavily doped contact layer and a second contact on thesubstrate; and i. slicing the substrate only in the trenches.